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Cite This: ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

Unusual Capacity Increases with Cycling for Ladder-Type Microporous Polymers Tyler B. Schon, So Young An, Andrew J. Tilley, and Dwight S. Seferos* Department of Chemistry, University of Toronto, 80 St. George Street, Toronto, Ontario M5S 2H6, Canada Department of Chemical Engineering and Applied Chemistry, University of Toronto, 200 College Street, Toronto, Ontario M5S 3E5, Canada

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S Supporting Information *

ABSTRACT: Microporous polymers using triptycene vertices and various ladder-type benzimidazole linkers are synthesized and tested as lithium-ion battery anodes. An unusual increase in performance is observed upon cycling, affording high capacities of 783 and 737 mAh g−1 for a perylene derivative and the pyromellitic derivative after 1000 cycles. The high performance of these materials after cycling is attributed to favorable electrode morphology and high crystallinity for perylene derivative, and the presence of charge carriers for pyromellitic derivative. By studying the effect of various linkers on the electrochemical performance, structure−property relationships are proposed that can be used to guide the development of high-performance materials for lithium-ion batteries. KEYWORDS: organic electrode materials, anodes, microporous polymers, lithium ion batteries, superlithiation, ladder-type polymer include using solid electrolytes with a high modulus,8,10 using highly concentrated electrolytes,9,11 and plating lithium metal within a carbon material.7,12 Although much progress has been made in this area, these strategies are not enough to prevent lithium dendrite formation after extended cycling and safety issues are still a major concern. Another way around this issue is to use low-voltage, high-capacity materials. Silicon,13−16 germanium,17,18 and phosphorus19 are examples of materials that have been widely investigated as new anode materials having theoretical capacities greater than 1000 mAh g−1 and voltages lower than 0.5 V vs Li/Li+. However, these materials experience a large expansion upon lithiation, sometimes greater than 300%, leading to low cycling stability. An alternative solution is to use organic materials that typically do not experience a large volume expansion.20 However, organic anode materials are relatively underdeveloped in part because of degradation at low potentials, especially compounds that have saturated carbons or complex functionality. Conjugated materials with a minimal amount of saturated carbons have been the focus of this field. One of the most popular classes of organic lithium-ion battery anodes are conjugated carboxylic acids.20−27 This is due to the low reduction potential of carboxyl groups (∼0.5 V vs Li/Li+) and the high abundance of carboxylic acidcontaining organic compounds. However, these compounds

Lithium-ion batteries are currently the best technology to address energy storage needs in electric vehicles, grid-scale storage, and consumer electronics. To enable the widespread adoption of these emerging applications, higher energy and power densities are needed, which requires extensive research into new electrode materials, electrode architectures, electrolytes, and cell design.1 Because the electrodes dictate the capacity of the battery, and therefore the energy density, much effort is devoted to developing materials that can store a large amount of charge on a per mass basis. Organic materials have emerged as promising candidates for high-capacity electrodes.2−5 However, the vast majority of organic electrodes for lithium-ion batteries are cathode materials. This is likely because the redox potential of most organic materials is more appropriate for the cathode and there is an obvious motivation for replacing the currently used metalbased cathodes due to the environmental burden, the unfavorable mechanical properties, and the high costs that are associated with them.6 However, the development of new rechargeable anode materials is important because the graphite electrodes that are used in commercial rechargeable lithiumion batteries have a relatively low capacity (∼300 mAh g−1) compared to lithium metal (∼3842 mAh g−1); and the formation of dendrites upon repeated cycling of lithium and sodium metal anodes poses serious safety concerns.7−9 Many studies have focused on addressing the inherent safety issues with alkali metal anodes by either inhibiting dendrite formation or replacing the anode with a new material altogether. Popular strategies to inhibit dendrite formation © XXXX American Chemical Society

Received: October 19, 2018 Accepted: January 3, 2019

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DOI: 10.1021/acsami.8b18293 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

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ACS Applied Materials & Interfaces Scheme 1. Synthesis of Fused Microporous Polymers in (a) Acidic and (b) Basic Conditionsa

a Note that in F-PDI-3-Tc, there will be a significant amount of singly reacted dianhydride in the product because of the higher ratio used in the reaction.

Figure 1. (a) FTIR spectra of synthesized microporous polymers, (b) XRD spectra of synthesized microporous polymers, (c) NMR spectra of FPDI-Tc, and (d) EPR spectra of synthesized microporous polymers.

still have a relatively low theoretical capacity (∼300 mAh g−1), which limits the energy density of the battery. Other classes of organic compounds have also been explored for lithium ion battery anodes such as imine-containing28,29 and carbonylcontaining compounds30−32 but these also suffer from low capacities (under 300 mAh g−1). Recently, an intriguing phenomenon has been discovered where the insertion of lithium ions into a conjugated structure can provide a capacity of more than 1500 mAh g−1.33 This phenomenon has been

termed “superlithiation” because of the large amount of lithium ions that are inserted into the conjugated organic structure. The proposed mechanism for this large capacity is a complete reduction of the sp2 framework where each carbon atom accepts one electron and one lithium ion (Scheme S1). This type of reduction has been observed in compounds such as aromatic and conjugated carboxylic acids34,35 and ladder-type polymers.36−38 Superlithiation compounds can greatly increase the energy density of the resulting batteries, however, there is a B


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ACS Applied Materials & Interfaces

398.58, 399.18, 398.98, and 399.08 eV in F-PDI-Tc, F-PDI-3Tc, F-NDI-Tc, and F-pyro-Tc corresponds to the sp2 nitrogen while the peaks at 400.68 eV for F-PDI-Tc, 400.88 eV for FPDI-3-Tc and F-NDI-Tc, and at 400.58 eV for F-pyro-Tc corresponds to the sp3 nitrogen. The results of the N 1s spectra show that the benzimidazole ring is formed between the dianhydride and the triptycene unit.4546 The XRD pattern of F-PDI-Tc has weak diffraction peaks that are similar to perylene diimide-based polymers linked through the diimide nitrogen, and also that of the perylene dianhydride starting material (Figure 1b).47 Peaks corresponding to 9.44, 7.21, 3.61, 3.25, and 3.21 Å come from the diffractions of the perylene units.48 The XRD pattern of F-PDI3-Tc is similar to that of F-PDI-Tc, but the peaks are much more intense and sharp indicating a higher crystallinity. Given that perylene π−π stacking interactions in the fully formed microporous polymer is expected to be frustrated by the bulky triptycene vertices, we attribute the higher signal intensity in the diffractogram of F-PDI-3-Tc to crystalline regions of singly reacted perylene dianhydride starting material leftover from the imidization reaction to form this material (as mentioned earlier). The F-NDI-Tc and F-pyro-Tc XRD patterns show broad amorphous halos with F-NDI-Tc having two peaks corresponding to a spacing of 6.14 and 3.79 Å while F-pyro-Tc has one very broad peak corresponding to a spacing of 4.29 Å. Because the arylene dianhydride starting materials are crystalline, this shows that F-NDI-Tc and F-pyro-Tc are functionalized to a relatively high degree which breaks up the crystal packing between the arylene units. These results suggest that the crystallinity of the microporous polymers is dictated by the strength of the interactions of the aromatic linker groups, with these interactions largely diminished in the more highly functionalized materials. Solid-state 13C NMR was performed to further characterize the materials (Figure 1c, Figure S3). The spectra of the microporous polymers are consistent with reports of similar compounds and the assignments correlate with those calculated by DFT.40 The chemical shift of the carbonyl carbon and the carbon adjacent to the two nitrogens in the benzimidazole ring are at 161, 159, 158, and 166 ppm, for FPDI-Tc, F-PDI-3-Tc, F-NDI-Tc, and F-pyro-Tc, respectively. Additionally, the sp3 carbon peak in the triptycene unit is observed at 54, 49, 52, and 53 ppm for F-PDI-Tc, F-PDI-3-Tc, F-NDI-Tc, and F-pyro-Tc, respectively, indicating that the incorporation of the triptycene unit into the materials is successful. Other peaks in the spectra correspond to the various aromatic carbons in the phenyl groups of the triptycene and the aromatic carbons in the arylene units. As the linker in the microporous polymers decreases in size from perylene to pyromellitic, the signal in the spectra decreases dramatically. Increasing the number of scans in the measurement did not provide greater resolution which suggests that the materials exhibit some paramagnetic character. Indeed, electron paramagnetic resonance (EPR) measurements show that all microporous polymers possess radical character with a small degree of anisotropy and hyperfine coupling for F-pyro-Tc and F-NDI-Tc and no hyperfine coupling for F-PDI-Tc and F-PDI3-Tc (Figure 1d). The g-factor is 2.0033, 2.0034, 2.0031, and 2.0030 for F-PDI-Tc, F-PDI-3-Tc, F-NDI-Tc, and F-pyro-Tc, respectively. The g-factors for the materials are similar to that observed for a similar ladder-type polymer, polybenzimidazobenzoisoquinoline (BBL), which has a g-factor of 2.0034.49,50 EPR studies of BBL show similar spectra to that obtained with

lack of systematic studies that relate the molecular structure of the material to the performance in a battery. Additionally, the use of microporous materials for superlithiation could lead to a high stability due to the permanent porosity, which relieves mechanical stress associated with ion influx.35−40 Here, we describe a series of ladder-type triptycene microporous polymers that incorporate benzimidazole linkers and their use as anode materials for lithium ion batteries. These compounds undergo superlithiation at low potentials and have an unusual cycling behavior providing a high capacity of 783 mAh g−1 for F-PDI-3-Tc after 1000 cycles. By studying the effect of linker on the performance of the battery, we establish important structure−property-function relationships for microporous polymers in lithium ion battery anodes that are capable of superlithiation. This work highlights the important discovery that the presence of charge-carriers in the material, in the form of radicals, can result in an extremely high performance even though the electrode morphology is far from ideal. The microporous polymers were synthesized in acidic or basic conditions using the corresponding dianhydride and the hexamino triptycene hydrochloride salt (1) or the free-base (2) respectively (Scheme 1). Using the hydrochloride salt of hexaamino triptycene (1) allows a gram scale preparation of 1 without the need for anhydrous or oxygen-free conditions compared to a synthesis using the highly air sensitive free-base 2.39 F-pyro-Tc and F-NDI-Tc were synthesized in 61.9% and quantitative yields, respectively, using the acidic conditions. However, attempts to synthesize F-PDI-Tc using the same acidic conditions only led to the recovery of starting material as indicated by the FTIR spectra (Figure S1). When F-PDI-Tc was synthesized in basic conditions, a new peak at 1690 cm−1 was observed that corresponds to the CO stretch and the CN stretch in the new 5-membered ring, which is consistent with similar small molecule organic compounds that contain the benzimidazole moiety (Figure 1a).36,40,41 We also synthesized an F-PDI-Tc material with a higher ratio of the dianhydride starting material, labeled F-PDI-3-Tc to determine the effect of crystallinity and ratio of perylene groups to triptycene linkers on the electrochemical performance. This material was, like the others, completely insoluble, suggesting that it is a polymeric material rather than a distinct small molecule. The FTIR of F-PDI-3-Tc was similar to that of FPDI-Tc, however, many of the peaks are very similar to that of the perylene dianhydride starting material suggesting that the degree of functionalization is low and/or there are defects in the material. When F-NDI-Tc and F-pyro-Tc were synthesized in acidic conditions, new peaks in the FTIR spectra emerged compared to their respective starting materials. Notably, the broad peaks at 1701, 1634, and 1549 cm−1 in F-NDI-Tc and 1721, 1625, and 1436 cm−1 in F-pyro-Tc are due to the C=O and C=N stretches (Figure 1a).36,40,41 The C 1s XPS spectra of the microporous polymers shows that all materials possess the correct carbon environments expected (Figure S4). Each material possesses a C−C carbon signal corresponding to the aliphatic carbon in the triptycene unit, a signal corresponding to the carbons in the triptycene rings bonded to nitrogen atoms, a signal corresponding to the benzimidazole carbons double bonded to oxygen and nitrogen, and a broad shakeup feature corresponding to the π−π* transition similar to that of other diimide containing compounds.42−44 Additionally, the N 1s XPS spectra shows two nitrogen bonding environments (Figure S5). The peaks at C


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ACS Applied Materials & Interfaces F-pyro-Tc and F-NDI-Tc, having a small degree of anisotropy and hyperfine coupling. The origin of the unpaired electrons in BBL was attributed to the formation of bond alternation charged defects that result in polarons, which was supported by an optical absorption in the near-infrared region (NIR) of the spectrum. The EPR spectra suggests that polarons are present in all of the synthesized materials, with the concentration of the polarons increasing from F-PDI-Tc to F-pyro-Tc. This increase in polarons is expected to be accompanied by an increase in conductivity. The conductivity of the materials was measured as pressed pellets in a two-point probe setup (Table S1). The conductivity for F-PDI-Tc, F-PDI-3-Tc, F-NDI-Tc, and F-pyro-Tc was 5.73 × 10−7 S cm−1, 5.65 × 10−10 S cm−1, 5.84 × 10−9 S cm−1, and 3.32 × 10−8 S cm−1 respectively. Although the trend in conductivity does not follow the same trend in apparent polaron concentration measured by EPR, grain boundaries and particle-to-particle contacts play a large role in the bulk conductivity of the materials and could be much different from the intrinsic conductivity of the materials.51 It is also noted that the quality of the F-NDI-Tc and F-pyro-Tc pellets was much poorer than the F-PDI-Tc and F-PDI-3-Tc pellets, even when pressed at approximately 400 MPa. The porosity and surface area of the materials were examined by CO2 gas adsorption at room temperature. All materials show a typical type I isotherm that is consistent with microporous materials (Figure S6). The surface area of the materials increases with decreasing aromatic linker size with FPDI-Tc, F-NDI-Tc, and F-pyro-Tc having surfaces areas of 222.0, 276.5, and 328.2 m2 g−1, respectively. The surface area of F-PDI-3-Tc was relatively low, only 124.0 m2 g−1, which is likely due to the close π−π stacking of the excess perylene units that decreases the pore volume. The pore size distribution calculated by DFT shows that all materials have similar pore sizes of 3.5, 5.0, and 7.9 Å for F-PDI-Tc 3.5, 4.8, and 8.2 Å for F-PDI-3-Tc, and 3.5, 5.0, and 8.2 Å for F-NDITc and F-pyro-Tc (Figure S7). To test the applicability of the materials for lithium-ion battery anodes, electrode films were cast as a composite of the active microporous polymer materials, carbon Super P, and a PVDF binder in a weight ratio of 60:30:10. The electrodes were first examined by SEM to characterize the morphology, which is important for the performance of the resultant lithium-ion batteries (Figure 2). The SEM images of F-PDI-Tc show that the active material is relatively well dispersed within the conductive carbon matrix in the electrode, with irregular shaped aggregates a few microns in size. The SEM images of the F-PDI-3-Tc electrode depicts a homogeneous distribution of the materials. The F-PDI-3-Tc material appears as small aggregates with an average length of less than 2 μm and a diameter of ∼300 nm that are distributed within the carbon Super P matrix. The F-NDI-Tc electrode has large aggregates distributed throughout the electrode of varying sizes with some less than a 3 μm in diameter and some up to 80 μm in length. However, there are some areas of the electrode that have a homogeneous distribution of the F-NDI-Tc material (Figure S8). The F-pyro-Tc electrode also has large aggregates in the electrode, although they are smaller in size than the F-NDI-Tc aggregates with lengths around 20 μm. The F-pyro-Tc electrode also has some regions where there is a homogeneous distribution of material. The large aggregate size in the F-NDI-Tc and F-pyro-Tc electrodes could have an effect on the electrochemical

Figure 2. SEM images of (a) F-PDI-Tc electrodes, (b) F-PDI-3-Tc electrodes, (c) F-NDI-Tc electrodes, and (d) F-pyro-Tc electrodes.

performance of the resultant lithium-ion battery due to the large lithium-ion diffusion length to penetrate the aggregate as well as the nonconductive nature of the microporous polymers.32 In order to decrease the aggregate size in the FNDI-Tc and F-pyro-Tc electrodes, and to examine its effect on lithium-ion battery performance, the materials and resultant slurry were subjected to ball milling. Examining the morphology of the resultant electrodes by SEM clearly shows a drastic reduction in the average aggregate size, and an obvious increase in electrode homogeneity (Figure S9). For the F-NDI-Tc electrode, ball-milling decreases the aggregate size to below 10 μm and an increase in homogeneity compared to the sample that was not ball-milled. The F-pyro-Tc electrode that was ball-milled also has a decreased aggregate size to diameters below 10 μm, there is a qualitative increase in the homogeneity of the film, and there is an apparent increase in number macropores in the film compared to the electrode cast without ball-milling. This large difference in film morphology for the F-pyro-Tc may suggest that there is chemical degradation occurring in the material upon ballmilling. Indeed, the FTIR spectra of the ball-milled F-NDI-Tc shows little change in the spectra while that of F-pyro-Tc changes significantly (Figure S10). This shows that F-NDI-Tc is stable to the ball-milling conditions while that of F-pyro-Tc is not stable. Therefore, only the effect of ball-milling on FNDI-Tc was examined, however this proved to have a negative effect on performance, as discussed in the Supporting Information. Lithium-ion battery coin cells were assembled using lithium as the reference and auxiliary electrode. From the cyclic voltammogram, the materials exhibit an electrochemical profile that is consistent with that observed for other superlithiation compounds (Figure S11).33,35,36,38,52 For the F-PDI-Tc electrode, the first scan has a cathodic peak at 2.23 V corresponding to the reduction of the carbonyl groups, and a reduction between 1 to 0 V with a cathodic peak at 0.33 V which is ascribed to the formation of the solid electrolyte interface. These peaks disappear in the consecutive scans while only a sloping cathodic current is observed in subsequent cycling. When reversing the scan, two small anodic peaks appear at 1.16 and 1.96 V, which correspond to the delithiation of the material. For the F-PDI-3-Tc electrode, the first scan reveals a sharp cathodic peak at 1.90 V which almost disappears in the consecutive scans. This peak is ascribed to D


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cycling, the electrode capacity was retested and showed an improvement. The capacity at 50 mA g−1 is 123 mAh g−1, which is almost five times the capacity before cycling. This increase in capacity is accompanied by a decrease in charge transfer resistance, as measure by impedance spectroscopy, from 385.7 to 72.91 Ω (Table S3). This suggests that cycling the electrode leads to an activation of the active material, as indicated by the lower resistance, and affords a higher capacity. The F-PDI-3-Tc electrode has an initial capacity of 95.7 mAh g−1 at a current density of 50 mA g−1. Increasing the current leads to a dramatic decrease in performance, with a capacity of only 15.3 mAh g−1 at a current density of 100 mA g−1. The charge profile of the F-PDI-3-Tc material has a plateau between 0.5 and 0.0 V and a long sloping discharge from 0 to 3 V, consistent with other reports of superlithiation. When the cycling stability of the F-PDI-3-Tc electrode is tested at 200 mA g−1, an almost linear increase in capacity starting a negligible contribution of the active material at the third cycle and increasing to 547 mAh g−1 on the 1000th cycle with an almost perfect Coulombic efficiency (Figure 3a). This increase in capacity upon cycling is much larger than that observed for F-PDI-Tc. After testing the cycling stability, the rate capabilities were re-examined and they show a marked improvement in performance (Figure 3b). At a current density of 50 mA g−1, the capacity of the cycled F-PDI-3-Tc electrodes reaches a maximum of 783 mAh g−1 and even at an extremely high current density of 1000 mA g−1, the electrode still maintains a capacity of 249 mAh g−1. The reason for this drastic rise in performance is likely due to an activation of the electrode by an increased penetration of the electrolyte within the crystal structure of F-PDI-3-Tc, leading to a decreased resistance. This resistance can be quantified by the impedance data, where the charge transfer resistance was found to from 214.7 Ω for the pristine electrode to 57.97 Ω for the cycled one (Figure 3c). Interestingly, post-mortem analysis of the electrode reveals that a significant amount of solid electrolyte interface (SEI) is formed, although the aggregate size of the FPDI-3-Tc material does not seem to change (Figure S19). It is, however, difficult to distinguish the different components of the electrode due to the thick SEI formed and the chemical similarities of the electrode components. F-NDI-Tc has a poor performance compared to F-PDI-Tc and F-PDI-3-Tc (Figure S14). Even at the lowest current density of 50 mA g−1, the material has a negligible contribution of the active material. This is likely due to the large aggregates in the electrode and poor crystallinity that inhibits efficient charge-transfer between distinct redox moieties. Cycling the electrode does not lead to an increased performance. The impedance spectra of the F-NDI-Tc electrode before and after cycling shows a small increase in the charge-transfer resistance from 380 Ω to 415.4 Ω. F-pyro-Tc has an initial performance similar to F-PDI-Tc and F-PDI-3-Tc (Figure S12). At a current density of 50 mA g−1, F-pyro-Tc has a capacity of 44.7 mAh g−1, and at a current density of 500 mA g−1, it has a capacity of 17.6 mAh g−1, indicating that it initially has a rate capability greater than that of F-PDI-Tc, F-PDI-3-Tc, and F-NDI-Tc before cycling. Cycling the battery leads to an increase in capacity from 30.7 mAh g−1 in the second cycle to 642 mAh g−1 in the 1000th cycle. Testing the rate capabilities after cycling results in a dramatic improvement in performance. The capacity at 50 mA g−1 is 737 mAh g−1 and at 1000 mA g−1 the capacity is 390 mAh g−1. The increase in capacity is also attributed to an

the reduction of the carbonyl oxygen and the reduction of the sp2 imine nitrogen.36 From 1 to 0 V vs Li/Li+, a reduction occurs that is attributed to the formation of the solid electrolyte interface and the insertion of lithium ions into the aromatic carbons in the material. When reversing the scan, there is a small, broad oxidation peak centered at 1.23 V that is ascribed to the deinsertion of lithium ions in the material. Beyond the first scan, the CV curves nearly overlap and there is a new anodic peak that emerges at 1.60 V. For F-NDI-Tc, the evolution of the profiles and the peaks are similar to F-PDI-Tc although there is much less current being drawn from the system. This may indicate that F-NDI-Tc is not as electrochemically active as F-PDI-Tc and F-PDI-3-Tc. The CV of Fpyro-Tc, however, is much different from the other two materials. The first scan contains cathodic peaks at 1.46, 0.92, and 0.42 V that correspond to the reduction of the carbonyl groups, the reduction of the imine groups, and the formation of the solid electrolyte interface, respectively. Additionally, from 1.0 to 0.0 V, there is a broad cathodic peak overlapping with the peaks at 0.92 and 0.42 V that correspond to the insertion of lithium ions into the aromatic carbons, much like that of the other materials. However, there is a much more pronounced anodic current upon reversing of the scan. Two distinct peaks at 0.53 and 1.01 V relate to the deinsertion of lithium ions. After the first scan, the peaks at 1.46 and 0.92 V are greatly diminished and the peak at 0.42 V completely disappears. The broad reduction and oxidation peaks from 0.0 to 1.0 V corresponding to the insertion and deinsertion of lithium ions into and out of the material become consistent after the first scan. These results show that redox activity is occurring at low potential and is consistent with what is observed for other superlithiation compounds. To determine the capacity of the materials, we performed galvanostatic charge−discharge (CD) experiments (performance metrics are summarized in Table 1). The F-PDI-Tc Table 1. Performance Metrics of the Microporous Polymer Batteries before and after Cyclinga sample F-PDITc F-PDI3-Tc F-NDITc F-pyroTc

capacity before cycling (mAh g−1)

capacity after 500 cycles (mAh g−1)

28.7

123

95.7

455 b

44.7

capacity after 1000 cycles (mAh g−1)

783

b

231

737

Capacities were measured at 50 mA g−1. bThe active material contribution to capacity is negligible to the overall capacity of the electrode. a

electrode has a capacity of 28.7 mAh g−1 at a current density of 50 mA g−1 (Figure S12). Increasing the current to 100 mA g−1 provides a low capacity of 0.9 mAh g−1. The charge profile of F-PDI-Tc has a plateau between 0.5 and 0.0 V and a long sloping discharge from 0 to 3 V, consistent with other reports of superlithiation.30−34 Cycling the electrode leads to a linear increase in capacity of from a negligible contribution of the active material on the second cycle to 71.7 mAh g−1 on the 500th cycle (Figure S13). This shows that there is a long activation period required for this material and prior to this activation the material suffers from poor kinetics.37 After E


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Figure 3. Cycling stability at 200 mA g−1 of (a) F-PDI-3-Tc and (d) F-pyro-Tc; CD curves of (b) F-PDI-3-Tc and (e) F-pyro-Tc after 1000 cycles; and impedance spectra of (c) F-PDI-3-Tc and (f) F-pyro-Tc at 0.1 V at different stages of cycling. The solid lines represent the fits of the impedance data.

synthesized for applications as lithium ion battery anodes. These materials undergo superlithiation when reduced at low potentials below 1 V vs Li/Li+. F-PDI-3-Tc has an initial capacity of 95.7 mAh g−1 but increases to 783 mAh g−1 after charging and discharging 1000 times. This unusual increase in capacity is attributed to an activation of the electrode that reduces the charge-transfer resistance, enhancing the kinetics of the redox reactions. The high performance of F-PDI-3-Tc is attributed to the high crystallinity giving rise to a short electron hopping distance between perylene units, a greater amount of redox-active perylene groups compared to F-PDI-Tc, and small aggregate size in the electrode that allows for short lithium-ion diffusion pathways. F-pyro-Tc shows similar performance, with an initial capacity of 44.7 mAh g−1 that increases to 737 mAh g−1 after 1000 cycles. The high performance of F-pyro-Tc is attributed to the larger presence of radicals in the material which provides a relatively high conductivity and a large pore volume that facilitates electrolyte penetration in the active material. The F-NDI-Tc electrode however showed poor performance due to the low crystallinity, the large aggregate size, and relatively low concentration of polarons. Reducing the aggregate size of F-NDI-Tc seemed to have a large effect on the electrochemical profile of the material, with more ideal battery behavior and well-defined voltage plateaus, but the film integrity was greatly diminished resulting in delamination from the current collector. Overall, it seems that ball-milling fails to improve the performance of F-NDI-Tc, evidenced by the impedance spectroscopy results showing similar charge-transfer resistances of the ball-milled sample and the sample that was not ball-milled. This work provides a guideline toward the design of organic superlithiation anodes and for electrode materials in general, proving that high performance can be attained without having favorable electrode morphology and crystallinity. Work into removing the need for an “activation” step of these electrode materials is expected to result in commercially viable materials.

activation of the electrode, similar to the F-PDI-3-Tc battery, however, this result is surprising due to the lack of crystallinity in F-pyro-Tc as well the relatively large aggregrate size in the electrode. The decrease in charge transfer resistance from 339.6 Ω before cycling to 30.05 Ω after cycling shows that the activation results in an increase in the charge transfer kinetics. Post mortem analysis of the F-pyro-Tc electrode reveals some SEI formation, although not as much as the F-PDI-3-Tc electrode, and almost no change in F-pyro-Tc aggregate size. The trend in performance of the microporous polymers can be explained by the crystallinity of the active material, the conductivity of the active material, and the morphology of the electrode. The excellent performance of F-PDI-3-Tc is explained by the favorable electrode morphology allowing for short ion diffusion pathways and short pathways to the conductive carbon material, the high crystallinity leading to a short electron hopping distance between the aromatic perylene units, and the high degree of redox activity afforded by the higher ratio of perylene groups. This allows F-PDI-3-Tc to have a high capacity after cycling, due to the high amount of active material accessed after electrode activation. The same mechanism explains the increased performance after cycling for the F-PDI-Tc electrode, although to a lesser extent due to the lower crystallinity and lower amount of perylene redox groups within the material. F-pyro-Tc has a relatively unfavorable morphology and a low crystallinity compared to that of F-PDITc and F-PDI-3-Tc. On the basis of previous studies for battery materials, it would be predicted that this compound would have a poor performance. However, its high performance can be attributed to the higher number of charge carriers within the material, indicated by the large EPR signal, which would result in a higher intrinsic conductivity. This is inferred from the higher bulk conductivity of F-pyro-Tc compared to FNDI-Tc, which has a similar crystallinity and particle size. The poor performance of F-NDI-Tc is attributed to the low crystallinity, poor morphology, and poor conductivity. In conclusion, novel ladder-type triptycene-based microporous polymers with benzimidazole linkers have been F


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of 150 mg mL−1 for F-PDI-Tc, F-PDI-3-Tc, and F-NDI-Tc and at a concentration of 75 mg−1 for F-pyro-Tc. The mixtures were then sonicated for 1 h, stirring every 15 min to homogenize. The slurries were then cast onto copper foils (McMaster-Carr) using a notch-bar with a height of 200 μm, according a previously published procedure.54 The electrodes were then dried in air on a hot plate at a temperature of 80 °C and then in a vacuum antechamber at a temperature of 65 °C before bringing into the glovebox. CR2023-type coin cells were purchased from MTI Corporation. A copper foil with a diameter of 16 mm (McMaster-Carr) was used as the anodic current collector, a lithium foil with a diameter of 16 mm was used as the reference/auxiliary electrode, and a Celgard polypropylene separator with a diameter of 19 mm was used to prevent short circuiting. An electrode punch (DPM Solution Inc.) was used to cut the electrodes to a 16 mm diameter and a hydraulic press (BT Innovations) was used to hermetically seal the cells. Approximately 80 μL of electrolyte (1:1 (v/v) ethylene carbonate:diethyl carbonate, 1 M LiPF6) was used to fill the cells prior to sealing. Reported capacities of the materials were calculated by subtracting the contribution from the Super P conductive additive; however, the graphs depict the capacity of the entire battery electrode.

EXPERIMENTAL METHODS

General Considerations. All reagents and electrolytes were purchased from Sigma-Aldrich and used as received. Organic solvents for synthetic procedures were obtained from Caledon Laboratories Ltd., dried in a solvent purification system (Innovative Technology) under argon and then further dried over 4 Å molecular sieves (SigmaAldrich). All electrochemical measurements and construction of lithium-ion batteries were performed at room temperature in an argon-filled glovebox (mBraun) with oxygen and moisture levels below 5 ppm. All battery capacities and current densities were calculated on the basis of the amount of the material in the electrode. Synthesis of 1 and 2. The synthesis of 1 and 2 was carried out using a previously reported procedure.39,53 The NMR spectra were consistent with what was reported. Synthesis of F-NDI-Tc and F-pyro-Tc. The synthesis of the fused arylene diimide materials was based on a procedure previously reported.36 Briefly, polyphosphoric acid (25 g) was added to an ovendried 3-necked flask fitted with an argon inlet. The viscous liquid was degassed by bubbling argon through it while heating at 110 °C for 24 h with stirring. The liquid was cooled to 50 °C and 1 (1 equiv.) was added and the solution was then stirred at 120 °C overnight. The solution was then cooled to 70 °C, and the corresponding dianhydride (1.5 equiv.) was added. The reaction was slowly heated to 180 °C at a heating rate of 4 °C min−1. Upon completion, the reaction was cooled and poured into 500 mL of rapidly stirring methanol. The residual solid in the reaction flask was dissolved in water and poured into the stirring methanol. After stirring for 20 min, the solid was filtered, dried, and ground with a mortar and pestle. The solid was then placed in a Soxhlet thimble and extracted with methanol for 24 h, followed by chloroform for 3 h. F-NDI-Tc: Performed on a 0.569 mmol scale (quantitative yield). 13C CP/MAS NMR δ: 157.86, 145.24, 137.03, 123.71, 51.73 ppm. Elemental analysis calculated for C41H14N6O3: C, 77.11; H, 2.21; N, 13.16. Found: C, 62.35; H, 3.01; N, 8.68. F-pyro-Tc: Performed on a 0.569 mmol scale (61.9% yield). 13C CP/MAS NMR δ: 165.95, 142.33, 132.48, 111.06, 53.30 ppm. Elemental anal. Calcd for C35H11N6O3: C, 74.60; H, 1.97; N, 14.91. Found: C, 55.39; H, 3.83; N, 10.38. Synthesis of F-PDI-Tc. A mixture of 2 (710 mg, 1.01 mmol), zinc acetate (148 mg, 0.808 mmol), perylene-3,4,9,10-tetracarboxylic acid dianhydride (594 mg, 1.515 mmol), and 27 g of imidazole was added to a flame-dried 3-necked flask fitted with a reflux condenser. The mixture was backfilled with argon three times to remove any oxygen. The reaction mixture was heated to 160 °C for 22 h. Upon completion, the reaction mixture was cooled slightly and then poured into a stirring solution of methanol. The solid was filtered through a Soxhlet thimble and was extracted with methanol for 1 day, acetone for 3 h, hexanes for 2.5 h, and chloroform for 18 h to yield a dark purple solid (82.1% yield). 13C CP/MAS NMR δ: 160.13, 142.12, 135.00, 126.46, 117.62, and 48.66 ppm. Elemental anal. Calcd for C92H32N6O12: C, 81.55; H, 2.44; N, 10.19. Found: C, 68.75; H, 3.50; N, 8.88. Synthesis of F-PDI-3-Tc. A mixture of 2 (163 mg, 0.473 mmol), zinc acetate (210 mg, 1.144 mmol), perylene-3,4,9,10-tetracarboxylic acid dianhydride (556 mg, 1.417 mmol), and 25 g of imidazole was added to a flame-dried 3-necked flask fitted with a reflux condenser. The mixture was backfilled with argon three times to remove any oxygen. The reaction mixture was heated to 160 °C for 22 h. Upon completion, the reaction mixture was cooled slightly and then poured into a stirring solution of methanol. The solid was filtered through a Soxhlet thimble and was extracted with methanol for 1 day, acetone for 3 h, hexanes for 2.5 h, and chloroform for 18 h to yield a dark purple solid (78.5% yield). 13C CP/MAS NMR δ: 160.13, 142.12, 135.00, 126.46, 117.62, 48.66 ppm. Elemental anal. Calcd for C92H32N6O12: C, 78.19; H, 2.28; N, 5.95. Found: C, 68.10; H, 2.58; N, 6.81. Preparation of Lithium-Ion Batteries. The materials were initially ground into a fine powder with a mortar and pestle. The ground materials were then mixed with carbon Super P and PVDF in a 60:30:10 (w/w/w) ratio and suspended in NMP at a concentration

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsami.8b18293.

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Additional experimental details, XPS, FTIR, gas adsorption, SEM images, NMR, TGA, and electrochemical properties (PDF)

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected] Twitter: @seferosgroup. ORCID

Dwight S. Seferos: 0000-0001-8742-8058 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by the Natural Sciences and Engineering Research Council of Canada (NSERC) and the University of Toronto Connaught Foundation. The authors thank Dr. Sergiy Nokhrin for help with EPR measurements, Prof. M. Taylor for the use of the FTIR spectrometer, and Prof. T. Bender for the use of the TGA. T.B.S. and S.Y.A. acknowledge support from NSERC.

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REFERENCES

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